Steel sheet and method of production of same

ABSTRACT

A steel sheet improved in hardenability and material formability having a predetermined chemical composition, characterized in that, in the metal structure of the steel sheet, an average grain size of carbides is 0.4 μm to 2.0 μm, an area ratio of pearlite is 6% or less, when a number of carbides in ferrite grains is A and a number of carbides at ferrite grain boundaries is B, B/A&gt;l, and when an X-ray diffraction intensity at {211}&lt;011&gt;at a plane of a part of ½ sheet thickness of the steel sheet is denoted by “I1” and an X-ray diffraction intensity at {100}&lt;011&gt;is denoted by “I0”, I1/I0&lt;1 is satisfied, and the steel sheet has a Vickers hardness of 100 HV to 150 HV.

TECHNICAL FIELD

The present invention relates to steel sheet and a method of productionof the same.

BACKGROUND ART

Steel sheet containing, by mass %, carbon in an amount of 0.1 to 0.7% isbeing used as a material for production of gears, clutches, and otherdrive system parts of automobiles by being used press-formed, enlargedin holes, bent, drawn, thickened, and thinned and cold forged bycombinations of the same from a blank. The strength of such parts issecured by quenching and tempering, so a high hardenability is demandedfrom steel sheet.

Furthermore, a high formability in the cold state is demanded from steelsheet used as a material for such drive system parts. Parts are mainlyformed by drawing and/or thickening. In forming parts, the biggestfactor governing the material formability is the plastic anisotropy.Improvement of the plastic anisotropy in steel sheet is necessary forapplication of steel sheet to the formation of parts.

Several proposals have been made up to now for the hardenabilitydemanded and formability improved in plastic anisotropy. The followingpatent literature discloses steel sheet excellent in cold forgeabilityand impact resistance characteristic.

For example, PLT 1 discloses, as steel for machine structural useimproving toughness by suppressing coarsening of crystal grains incarburization heat treatment, steel for machine structural usecontaining, by mass %, C: 0.10 to 0.30%, Si: 0.05 to 2.0%, Mn: 0.10 to0.50%, P: 0.030% or less, S: 0.030% or less, Cr: 1.80 to 3.00%, Al:0.005 to 0.050%, Nb: 0.02 to 0.10%, and N: 0.0300% or less and having abalance of Fe and unavoidable impurities, having a structure before coldworking comprised of ferrite and pearlite structures, and having anaverage value of ferrite grain size of 15 μm or more.

PLT 2 discloses, as steel excellent in cold workability and carburizingand quenching ability, steel containing C: 0.15 to 0.40%, Si: 1.00% orless, Mn: 0.40% or less, sol. Al: 0.02% or less, N: 0.006% or less, andB: 0.005 to 0.050%, having a balance of Fe and unavoidable impurities,and having a structure mainly comprised of ferrite phases and graphitephases.

PLT 3 discloses a steel material for carburized bevel gear use excellentin impact strength, a high toughness carburized bevel gear, and a methodof production of the same.

PLT 4 discloses, for a part produced by spheroidal annealing, then acold forging and a carburizing, quenching, and tempering process, steelfor carburized part use having excellent workability while suppressingcoarsening of crystal grains even with subsequent carburization andhaving an excellent impact resistance characteristic and impact fatigueresistance characteristic.

PLT 5 discloses as cold tool steel for plasma carburization use a steelcontaining C: 0.40 to 0.80%, Si: 0.05 to 1.50%, Mn: 0.05 to 1.50%, andV: 1.8 to 6.0%, further containing one or more of Ni: 0.10 to 2.50%, Cr:0.1 to 2.0%, and Mo: 3.0% or less, and having a balance of Fe andunavoidable impurities.

On the other hand, there have been the following proposals forimprovement of the formability, that is, the improvement of plasticanisotropy.

For example, PLT 6 proposes prescribing the carbide grain size andspheroidization rate in steel containing C: 0.25 to 0.75% and improvingthe in-plane anisotropy by the cold rolling rate and box annealingconditions, the coiling temperature in hot rolling, and provisions onthe texture so as to limit the “r” value and Δr.

PLTs 7 and 8 propose to prescribe the heating and annealing conditionsof a hot rolled material between stands of a finish rolling machine soas to reduce the Δr value and improve the in-plane anisotropy. PLT 8proposes steel sheet reduced in in-plane anisotropy by prescribing hotrolling during which performing finish rolling at a temperature of theAr3 point or more and coiling at 500 to 630° C.

CITATION LIST Patent Literature

PLT 1: Japanese Patent Publication No. 2013-040376A

PLT 2: Japanese Patent Publication No. 06-116679A

PLT 3: Japanese Patent Publication No. 09-201644A

PLT 4: Japanese Patent Publication No. 2006-213951A

PLT 5: Japanese Patent Publication No. 10-158780A

PLT 6: Japanese Patent Publication No. 2000-328172A

PLT 7: Japanese Patent Publication No. 2001-073076A

PLT 8: Japanese Patent Publication No. 2001-073077A

SUMMARY OF INVENTION Technical Problem

The above patent literature proposed improvement of the in-planeanisotropy, but did not propose the provision of the strength demandedfrom the part, that is, the hardenability.

The present invention was made in consideration of the above situationin the prior art and has as its object the provision of steel sheetimproved in hardenability and material formability, in particular,optimal for obtaining a gear or other part by thickening or other coldforging, and a method of production of the same.

Solution to Problem

To solve the above problem and obtain steel sheet suitable for thematerial of a drive system part etc., it can be understood that in steelsheet containing the C necessary for raising the hardenability, it issufficient to increase the grain size of the ferrite, spheroidize thecarbides (mainly cementite) by a suitable grain size, and decrease thepearlite structures. This is due to the following reasons.

Ferrite phases are low in hardness and high in ductility. Therefore, ina structure mainly comprised of ferrite, it becomes possible to increasethe grain size so as to raise the material formability.

Carbides, by being made to suitably disperse in the metal structure, canmaintain the material formability while imparting an excellent wearresistance and rolling fatigue characteristic, so are structuresessential for drive system parts. Further, the carbides in the steelsheet are strong particles obstructing slip. By forming carbides at theferrite grain boundaries, it is possible to prevent propagation of slipexceeding the crystal grain boundaries and suppress the formation ofshear zones. The cold forgeability is improved and, simultaneously, theformability of steel sheet is also improved.

However, cementite is a hard, brittle structure. If a laminar structurewith ferrite present, that is, in the state of pearlite, the steelbecomes hard and brittle, so it has to be present in a spheroidal form.If considering the cold forgeability and the occurrence of fractures atthe time of forging, its grain size has to be a suitable range.

However, no method of production for realizing the above structure hasbeen disclosed up to now. Therefore, the inventors intensivelyresearched a method of production for realizing the above structure.

As a result, they discovered the following: To make the metal structureof the steel sheet after coiling after hot rolling a bainite structureof fine pearlite or fine ferrite with small lamellar spacing in whichcementite is dispersed, the steel sheet is coiled at a relatively lowtemperature (400° C. to 550° C.). By coiling at a relatively lowtemperature, the cementite dispersed in the ferrite also easily becomesspheroidal. Next, the cementite is partially made spheroidal byannealing at a temperature just under the Ac1 point as first stageannealing. Next, as second stage annealing, part of the ferrite grainsis left while part is transformed to austenite by annealing at atemperature between the Ac1 point and Ac3 point (so-called dual phaseregion of ferrite and austenite). By then making the remaining ferritegrains grow while slowly cooling the steel while using these as nucleito transform the austenite to ferrite, it is possible to obtain largeferrite phases and make cementite precipitate at the grain boundaries torealize the above structure.

That is, the inventors found that it is difficult to realize a method ofproduction of steel sheet satisfying both hardenability and formabilityeven if adjusting the heat rolling conditions, annealing conditions,etc. separately and that it is possible to realize this by optimizationby a so-called integrated process of hot rolling, annealing, etc.

Further, they found that for improvement of the drawability at the timeof cold forming, reduction of the plastic anisotropy is necessary and,to improve this, adjustment of the hot rolling conditions is important.

The present invention was made based on these findings and has as itsgist the following:

(1) A steel sheet consisting of, by mass %, C: 0.10 to 0.70%, Si: 0.01to 0.30%, Mn: 0.30 to 3.00%, Al: 0.001 to 0.10%, Cr: 0.010 to 0.50%, Mo:0.0010 to 0.50%, B: 0.0004 to 0.01%, Ti: 0.001 to 0.10%, P: 0.02% orless,

S: 0.01% or less, N: 0.0200% or less, 0: 0.0200% or less, Sn: 0.05% orless, Sb: 0.05% or less, As: 0.05% or less, Nb: 0.10% or less, V: 0.10%or less, Cu: 0.10% or less, W: 0.10% or less, Ta: 0.10% or less, Ni:0.10% or less, Mg: 0.05% or less, Ca: 0.05% or less, Y: 0.05% or less,Zr: 0.05% or less, La: 0.05% or less, and Ce: 0.05% or less and abalance of Fe and unavoidable impurities, wherein the metal structure ofthe steel sheet includes carbide having an average grain size of 0.4 μmto 2.0 μm, perlite having an area ratio of 6% or less, and ferritewherein a ratio of a number of carbides at ferrite grain boundaries to anumber of carbides in ferrite grains of over 1; and I1/I0<1 beingsatisfied when an X-ray diffraction intensity at {211}<011>at a plane ofa part of ½ sheet thickness of the steel sheet is denoted by “I1” and anX-ray diffraction intensity at {100}<011>is denoted by “I0”, the steelsheet having a Vickers hardness of 100 HV to 150 HV.

(2) A method of production for producing steel sheet according to (1)comprising hot rolling a steel slab of a chemical composition accordingto (1) with finish rolling temperature between 820° C. and 950° C., toobtain hot rolled steel sheet; coiling the hot rolled steel sheet at400° C. to 550° C.; pickling the coiled hot rolled steel sheet; heatingthe pickled hot rolled steel sheet to an annealing temperature of 650°C. to 720° C. by a heating rate of 30° C/hour to 150° C/hour and holdingthe steel sheet for 3 hours to 60 hours as a first stage of annealing;next, heating the hot rolled steel sheet to an annealing temperature of725° C. to 790° C. by a heating rate of 1° C/hour to 80° C/hour andholding the steel sheet for 3 hours to less than 10 hours as a secondstage of annealing; and, next, cooling the annealed hot rolled steelsheet to 650° C. by a cooling rate of 1° C/hour to 100° C/hour.

Advantageous Effects of Invention

According to the present invention, it is possible to provide steelsheet excellent in hardenability and material formability, inparticular, optimal for obtaining a gear or other part by forming bythickening or other cold forging, and a method of production of thesame.

DESCRIPTION OF EMBODIMENTS

Below, the present invention will be explained in detail. First, thereasons for limitation of the chemical composition of the steel sheet ofthe present invention will be explained. Here, the “%” according to thechemical composition means “mass %”.

C: 0.10 to 0.70%

C is an element forming carbides and effective for strengthening thesteel and refining the ferrite grains. To suppress the formation of amatte surface in cold working and secure surface beauty of a cold forgedpart, suppression of coarsening of the ferrite grain size is necessary.

If C is less than 0.10%, the carbides become insufficient in volumefraction and coarsening of the carbides during annealing can no longerbe suppressed, so C is made 0.10% or more. Preferably it is 0.14% ormore. On the other hand, if the content of C increases, the carbidesincrease in volume fraction, cracks are formed acting as starting pointsof breakage at the time of an instantaneous load, and there is the fearthat the formability and impact resistance characteristic will fall. Ifmaking this drop as small as possible, C is made 0.40% or less.Preferably it is 0.38% or less.

On the other hand, if the carbides increase in volume fraction and thestrength rises, the fatigue characteristic is improved, so whenimproving the fatigue characteristic, C is made over 0.40%. Preferablyit is 0.44% or more. If C is over 0.70%, a large amount of cracksforming starting points of breakage are formed and the fatiguecharacteristic conversely falls, so C is made 0.70% or less. Preferablyit is 0.66% or less.

Si: 0.01 to 0.30%

Si is an element which acts as a deoxidizing agent and further has aneffect on the form of the carbides and contributes to the improvement ofthe material formability. To obtain the deoxidizing effect, Si is made0.01% or more. Preferably it is 0.07% or more.

If Si is over 0.30%, due to solution strengthening of the ferrite, thehardness rises and the ductility falls, fractures easily occur at thetime of cold forging, and the formability at the time of cold forgingand the impact resistance characteristic after carburization, quenching,and temperature falls, so Si is made 0.30% or less. Preferably it is0.28% or less.

Mn: 0.30 to 3.00%

Mn is an element controlling the form of carbides in two-stageannealing. If less than 0.30%, in the gradual cooling after second stageannealing, it becomes difficult to form carbides at the ferrite grainboundaries, so Mn is made 0.30% or more. Preferably it is 0.40% or more.

If Mn is over 1.00%, after carburization, quenching, and tempering, thetoughness falls, but on the other hand, the strength rises. When tryingto keep down the drop in toughness after carburization, quenching, andtempering as much as possible Mn is made 1.00% or less. Preferably it is0.96% or less.

When trying to raise the strength, Mn is made over 1.00%. Preferably itis 1.10% or more. If Mn is over 3.00%, after carburization, quenching,and tempering, the toughness remarkably falls, so Mn is made 3.00% orless. Preferably it is 2.70% or less.

Al: 0.001 to 0.10%

Al is an element which acts as a deoxidizing agent and stabilizesferrite. If less than 0.001%, the effect of addition is not sufficientlyobtained, so Al is made 0.001% or more. Preferably it is 0.004% or more.

On the other hand, if Al is over 0.10%, the number of carbides at theferrite grain boundaries decreases and the formability falls, so Al ismade 0.10% or less. Preferably it is 0.09% or less.

Cr: 0.010 to 0.50%

Cr is an element effective for stabilization of carbides at the time ofheat treatment. If less than 0.010%, it becomes difficult to causecarbides to remain at the time of carburization, coarsening of theaustenite grain size at the surface layer is invited, and the strengthdrops, so Cr is made 0.010% or more. Preferably it is 0.050% or more.

On the other hand, if Cr is over 0.50%, the amount of Cr concentratingat the carbides increases and a large amount of fine carbides remain inthe austenite phases produced by the two-stage annealing, carbidesremain in the ferrite grains after gradual cooling inviting an increasein the hardness and a decrease in the number of carbides at the ferritegrain boundaries fall, and the formability falls, so Cr is made 0.50% orless. Preferably it is 0.40% or less.

Mo: 0.001 to 0.50%

Mo, like Mn and Cr, is an element effective for control of the form ofcarbides. If less than 0.001%, the effect of addition is not obtained,so Mo is made 0.001% or more. Preferably it is 0.005% or more.

On the other hand, if over 0.50%, Mo concentrates at the carbides,stable carbides increase even in the austenite phases, carbides remaininside the ferrite grains after gradual cooling inviting an increase inthe hardness and a decrease in the number of carbides at the ferritegrain boundaries, and the material formability falls, so Mo is made0.50% or less. Preferably it is 0.40% or less.

B: 0.0004 to 0.01%

B is an element raising the hardenability and further raising thetoughness. In the steel sheet of the present invention, a predeterminedhardenability is required, so 0.0004 to 0.01% is added. If less than0.0004%, the effect of addition is not obtained, so B is made 0.0004% ormore. Preferably it is 0.0010% or more.

On the other hand, if over 0.01%, coarse B compounds becoming the causeof internal defects and other flaws at the time of steel production areformed, so B is made 0.01% or less. Preferably it is 0.007% or less.

Ti: 0.001 to 0.10%

Ti is an element forming nitrides and contributing to refinement of thecrystal grains and works to effectively bring out the effect of additionof B. If less than 0.001%, the effect of addition is not obtained, so Tiis made 0.001% or more. Preferably it is 0.010% or more.

On the other hand, if over 0.10%, coarse Ti nitrides are formed and thematerial formability falls, so Ti is made 0.10% or less. Preferably itis 0.07% or less.

The following elements are impurities and have to be controlled tocertain amounts or less.

P: 0.02% or less

P is an element segregating at the ferrite grain boundaries and workingto suppress the formation of carbides at the ferrite grain boundaries.For this reason, the smaller the amount of P, the better. The content ofP may also be 0, but if reducing it to less than 0.0001%, the refiningcosts greatly increase, so the substantive lower limit is 0.0001 to0.0013%.

If P is over 0.02%, formation of carbides at the ferrite grainboundaries is suppressed, the number of carbides decreases, and thematerial formability falls, so P is made 0.02% or less. Preferably it is0.01% or less.

S: 0.01% or less

S is an impurity element forming MnS and other nonmetallic inclusions.The nonmetallic inclusions form starting points of fracture at the timeof cold forging, so the smaller the S, the better. The content of S mayalso be 0, but to lower S to less than 0.0001%, the refining costsgreatly increase, so the substantive lower limit is 0.0001 to 0.0012%.

If S is over 0.01%, nonmetallic inclusions are formed and the materialformability falls, so S is made 0.01% or less. Preferably it is 0.009%or less.

N: 0.02% or less

N is an element which, if present in a large amount, causesembrittlement of the ferrite. For this reason, the smaller the amount ofN, the better. The content of N may also be 0, but to lower N to lessthan 0.0001%, the refining costs greatly increase, so the substantivelower limit is 0.0001 to 0.0006%.

If N is over 0.02%, the ferrite becomes brittle and the materialformability falls, so N is made 0.02% or less. Preferably it is 0.017%or less.

When the steel sheet of the present invention contains C: 0.10 to 0.40%and Mn: 0.30 to 1.00%, embrittlement of the ferrite is suppressed, so Nis made 0.01% or less. Preferably it is 0.007% or less.

O: 0.02% or less O is an element which, if present in a large amount,promotes the formation of coarse oxides. For this reason, the smallerthe amount of O, the better, but to lower O to less than 0.0001%, therefining costs greatly increase, so the amount is made 0.0001% or more.Preferably it is 0.0011% or more.

On the other hand, if over 0.020%, coarse oxides are formed in thesteel, the oxides become starting points of fracture at the time of coldforging, and the material formability falls, so O is made 0.02% or less.Preferably it is 0.01% or less.

Sn: 0.05% or less

Sn is an element which unavoidably enters from the steel startingmaterials. For this reason, the smaller the amount of Sn, the better.The content of S may also be 0, but to lower S to less than 0.001%, therefining costs greatly increase, so the substantive lower limit is 0.001to 0.002%.

On the other hand, if over 0.05%, the ferrite becomes brittle and thematerial formability falls, so Sn is made 0.05% or less. Preferably itis 0.04% or less.

Sb: 0.05% or less

Sb, like Sn, is an element which unavoidably enters from the steelstarting materials, segregates at the ferrite grain boundaries, andreduces the number of carbides at the ferrite grain boundaries. For thisreason, the smaller the amount of Sb, the better. The content of Sb mayalso be 0, but to lower Sb to less than 0.001%, the refining costsgreatly increase, so the substantive lower limit is 0.001 to 0.002%.

On the other hand, if over 0.050%, Sb segregates at the ferrite grainboundaries, the number of carbides at the ferrite grain boundariesdecreases, and the material formability falls, so Sb is made 0.050% orless. Preferably it is 0.04% or less.

As: 0.05% or less

As, like Sn and Sb, is an element which unavoidably enters from thesteel starting materials and segregates at the ferrite grain boundaries.For this reason, the smaller the amount of As, the better. The contentof As may also be 0, but to lower As to less than 0.001%, the refiningcosts greatly increase, so the substantive lower limit is 0.001 to0.002%.

On the other hand, if over 0.05%, As segregates at the ferrite grainboundaries, the number of carbides at the ferrite grain boundariesdecreases, and the material formability falls, so As is made 50% orless. Preferably it is 0.04% or less.

The steel sheet of the present invention has the above elements as basicelements, but may further contain the following elements for the purposeof improving the cold forgeability of the steel sheet. The followingelements are not essential for obtaining the effects of the presentinvention, so the contents may also be 0.

Nb: 0.10% or less

Nb is an element effective for control of the form of the carbides.Further, it is an element refining the structure and contributing toimprovement of the toughness. To obtain this effect of addition, Nbpreferably is made 0.001% or more. More preferably it is 0.002% or more.

On the other hand, if over 0.10%, a large number of fine Nb carbidesprecipitate, the strength excessively rises, and, further, the number ofcarbides at the grain boundaries falls and the cold forgeability falls,so Nb is made 0.10% or less. Preferably it is 0.09% or less.

V: 0.10% or less

V, like Nb, is an element effective for control of the form of thecarbides. Further, it is an element refining the structure andcontributing to improvement of the toughness. To obtain this effect ofaddition, V preferably is made 0.01% or more. More preferably it is0.004% or more.

On the other hand, if over 0.10%, a large number of fine V carbides areformed, the strength rises too much, the number of carbides at theferrite grain boundaries decreases, and the material formability falls,so V is made 0.10% or less. Preferably it is 0.09% or less.

Cu: 0.10% or less

Cu is an element segregating at the ferrite grain boundaries. Further,it is an element forming fine precipitates and contributing to theimprovement of strength. To obtain the effect of improvement ofstrength, Cu preferably is made 0.001% or more. More preferably it is0.008% or more.

On the other hand, if over 0.10%, segregation at the ferrite grainboundaries invites red shortness and causes the productivity in hotrolling to fall, so Cu is made 0.10% or less. Preferably it is 0.09% orless.

W: 0.10% or less

W, like Nb and V, is an element effective for control of the form ofcarbides. To obtain this effect of addition,

W preferably is made 0.001% or more. More preferably it is 0.003% ormore.

On the other hand, if over 0.10%, a large number of fine W carbides areformed, the strength rises too much, the number of carbides at theferrite grain boundaries decreases, and the material formability falls,so W is made 0.10% or less. Preferably it is 0.08% or less.

Ta: 0.001 to 0.10%

Ta, like Nb, V, and W, is an element effective for control of the formof carbides. To obtain this effect of addition, Ta preferably is made0.001% or more. More preferably it is 0.007% or more.

On the other hand, if over 0.10%, a large number of fine Ta carbides areformed, the strength rises too much, the number of carbides at theferrite grain boundaries decreases, and the material formability falls,so T is made 0.100% or less. Preferably it is 0.09% or less.

Ni: 0.10% or less

Ni is an element effective for improvement of the impact resistancecharacteristic of the formed part. To obtain this effect of addition, Nipreferably is made 0.001% or more. More preferably it is 0.002% or more.

On the other hand, if over 0.10%, the number of carbides at the ferritegrain boundaries decreases and the material formability falls, so Ni ismade 0.10% or less. Preferably it is 0.09% or less.

Mg: 0.05% or less

Mg is an element which can control the form of sulfides by addition in atrace amount. To obtain this effect of addition, Mg preferably is made0.0001% or more. More preferably it is 0.0008% or more.

On the other hand, if over 0.05%, the ferrite becomes brittle and thematerial formability falls, so Mg is made 0.05% or less. Preferably itis 0.04% or less.

Ca: 0.05% or less

Ca, like Mg, is an element which can control the form of sulfides byaddition in a trace amount. To obtain this effect of addition, Capreferably is made 0.001% or more. More preferably it is 0.003% or more.

On the other hand, if over 0.05%, coarse Ca oxides are formed and becomestarting points of fracture at the time of forming by cold forging, thatis, the material formability falls, so Ca is made 0.05% or less.Preferably it is 0.04% or less.

Y: 0.05% or less

Y, like Mg and Ca, is an element which can control the form of sulfidesby addition in a trace amount. To obtain this effect of addition, Ypreferably is made 0.001% or more. More preferably it is 0.003% or more.

On the other hand, if over 0.05%, coarse Y oxides are formed and becomestarting points of fracture at the time of forming by cold forging, thatis, the material formability falls, so Y is made 0.05% or less.Preferably it is 0.03% or less.

Zr: 0.05% or less

Zr, like Mg, Ca, and Y, is an element which can control the form ofsulfides by addition in a trace amount. To obtain this effect ofaddition, Zr preferably is made 0.001% or more. More preferably it is0.004% or more.

On the other hand, if over 0.05%, coarse Zr oxides are formed and becomestarting points of fracture at the time of forming by cold forging, thatis, the material formability falls, so Zr is made 0.05% or less.Preferably it is 0.04% or less.

La: 0.05% or less

La is an element able to control the form of the sulfides by addition ina trace amount, but is an element which segregates at the grainboundaries and reduces the number of carbides at the ferrite grainboundaries. To obtain the effect of control of the form of sulfides, Lapreferably is made 0.001% or more. More preferably it is 0.003% or more.

On the other hand, if over 0.05%, La segregates at the ferrite grainboundaries, the number of carbides at the ferrite grain boundariesdecreases, and the material formability falls, so La is made 0.05% orless. Preferably it is 0.04% or less.

Ce: 0.05% or less

Ce, like La, is an element able to control the form of the sulfides byaddition in a trace amount, but is an element which segregates at thegrain boundaries and reduces the number of carbides at the ferrite grainboundaries. To obtain the effect of control of the form of sulfides, Cepreferably is made 0.001% or more. More preferably it is 0.003% or more.

On the other hand, if over 0.05%, Ce segregates at the ferrite grainboundaries, the number of carbides at the ferrite grain boundariesdecreases, and the material formability falls, so Ce is made 0.050% orless. Preferably it is 0.04% or less.

The balance of the chemical composition is Fe and unavoidableimpurities.

Next, the structure of the steel sheet of the present invention will beexplained.

The structure of the steel sheet of the present invention is a structuresubstantially comprised of ferrite and carbides. Carbides are compoundsof iron and carbon of cementite (Fe₃C) plus compounds of cementite inwhich Fe atoms are substituted by Mn, Cr, and other alloy elements andalloy carbides (M₂₃C₆, M₆C, MC, etc. [M: Fe and other metal elementsadded as alloys]).

When forming steel sheet into a predetermined part shape, a shear zoneis formed at the macrostructure of the steel sheet and slip deformationoccurs concentrated near the shear zone. In slip deformation, along withproliferation of dislocations, a region of a high dislocation density isformed near the shear zone. Along with the increase in the amount ofstrain imparted to the steel sheet, slip deformation is promoted and thedislocation density increases.

In cold forging, strong working is performed with an equivalent strainexceeding 1. For this reason, in conventional steel sheet, it was notpossible to prevent the formation of voids and/or cracks along with theincrease in dislocation density and was difficult to improve the coldforgeability. To solve this problem, it is effective to suppress theformation of a shear zone at the time of forming.

From the viewpoint of the microstructure, formation of a shear zone canbe understood as the phenomenon of slip occurring at a certain one graincrossing the crystal grain boundary and being continuously propagated tothe adjoining grain. Accordingly, to suppress the formation of a shearzone, it is necessary to prevent propagation of slip crossing crystalgrain boundaries.

The carbides in steel sheet are strong particles inhibiting slip. Byforming carbides at the ferrite grain boundaries, it becomes possible toprevent the propagation of slip crossing crystal grain boundaries andsuppress the formation of a shear zone and improve the coldforgeability. Simultaneously, the steel sheet is also improved informability.

The formability of steel sheet is largely due to the accumulation ofstrain inside the crystal grains (accumulation of dislocations). Ifpropagation of strain to the adjoining crystal grains is blocked at thecrystal grain boundaries, the amount of strain inside the crystal grainsincreases. As a result, the work hardening rate increases and theformability is improved.

To obtain such an effect, carbides have to be made to disperse in themetal structure in suitable sizes. Therefore, the average grain size ofcarbides is made 0.4 μm to 2.0 μm. If the average grain size of thecarbides is less than 0.4 μm, the steel sheet remarkably increases inhardness and falls in cold forgeability.

More preferably it is 0.6 μm or more.

On the other hand, if the average particle size of the carbides exceeds2.0 μm, at the time of cold forming, the carbides form starting pointsof cracks. More preferably, it is 1.95 μm or less.

Further, cementite, a carbide of iron, is a hard and brittle structure.If present in the form of pearlite, which is a layered structure withferrite, the steel becomes hard and brittle. Therefore, pearlite has tobe reduced as much as possible. In the steel sheet of the presentinvention, the area ratio is made 6% or less.

Pearlite has a unique lamellar structure, so can be discerned byobservation by an SEM or optical microscope. By calculating the regionof the lamellar structure at any cross-section, the area ratio of thepearlite can be found.

Based on theory and principle, cold forgeability is considered to bestrongly affected by the rate of coverage of the ferrite grainboundaries by carbides. High precision measurement is sought, butmeasurement of the rate of coverage of ferrite grain boundaries bycarbides in a three-dimensional space requires serial sectioning SEMobservation using an FIB to repeatedly cut a sample for observation in ascanning electron microscope or 3D EBSP observation. A massivemeasurement time is required and technical knowhow has to be built up.

The inventors judged that the above method of observation was not ageneral method of analysis and did not employ it. They searched for asimpler, higher precision indicator for evaluation. As a result, theydiscovered that it is possible to quantitatively evaluate the coldforgeability and formability by using the ratio B/A of the number B ofcarbides at the ferrite grain boundaries to the number A of carbides inthe ferrite grains as an indicator and that if the ratio B/A exceeds 1,the cold forgeability and the formability in drawing and thickeningremarkably rise.

Buckling, folding, and twisting of the steel sheet occurring at the timeof cold forging occur due to localization of strain accompanying theformation of a shear zone, so by forming carbides at the ferrite grainboundaries, the formation of a shear zone and localization of strain arereduced and buckling, folding, and twisting are suppressed.

The carbides are observed by a scanning electron microscope. Beforeobservation, the sample for observation of the structure is polished bywet polishing by Emery paper and a diamond abrasive having an averageparticle size of 1 μm, the observed surface is polished to a mirrorfinish, then a 3% nitric acid-alcohol solution is used to etch thestructure. The magnification of the observation was made 3000× andimages of eight fields of 30 μm×40 μm at a sheet thickness ¼ layer werecaptured at random.

The obtained structural images were analyzed by image analyzing software(Win ROOF made by Mitani Shoji) to measure in detail the areas of thecarbides contained in the analyzed regions. The circle equivalentdiameters (=2×√(area/3.14)) were found from the areas of the carbidesand the average value was made the grain size of the carbides. Notethat, to keep down the effect of measurement error due to noise,carbides with an area of 0.01 μm² or less are excluded from the coverageof the evaluation.

The number of carbides present at the ferrite grain boundaries arecounted, the number of carbides at the grain boundaries are subtractedfrom the total number of carbides, and the number of carbides in theferrite grains are found. Based on the measured and calculated number ofcarbides, the ratio B/A of the number B of carbides at the ferrite grainboundaries with respect to the number A of carbides inside the ferritegrains is calculated.

In the structure of the steel sheet after annealing, the ferrite grainsize is preferably 3 μm to 50 μm from the viewpoint of improvement ofthe cold forgeability. If the ferrite grain size is less than 3 μm, thehardness increases and fractures and cracks easily form at the time ofcold forging, so the ferrite grain size is preferably 3 μm or more. Morepreferably it is 5 μm or more.

If the ferrite grain size is over 50.0 μm, the number of carbides on thecrystal grain boundaries suppressing the propagation of slip isdecreased and the cold forgeability falls, so the ferrite grain size ispreferably 50 μm or less. More preferably it is 40 μm or less.

The ferrite grain size is measured by using the above-mentionedprocedure to polish the observed surface of the sample surface to amirror finish, then etching it by a 3% nitric acid-alcohol solution andobserving the structure by an optical microscope or scanning electronmicroscope and applying the line segment method to the captured image.

At the time of cold forging, in addition to control of the form ofcarbides, drawability at the time of cold forging becomes necessary.

To improve the drawability at the time of cold forging, improvement ofthe plastic anisotropy becomes necessary. For this reason, control ofthe texture at the hot rolled steel sheet is necessary. The texture isevaluated by X-ray diffraction at a plane parallel to the sheet surfaceat a ½ sheet thickness part of the hot rolled steel sheet. For X-raydiffraction, X-rays from a Mo tube are used.

The diffraction intensities at the diffraction orientations {110},{220}, {211}, and {310} due to reflection are obtained and based onthese an ODF is prepared. For the preparation of the ODF, thediffraction intensity data of random orientations of iron is used. Fromthis, the X-ray diffraction intensity of {211}<011>is found as I1 andthe X-ray diffraction intensity of {100}<011>is found as I0. If thisI1/I0 is less than 1, it means that the recrystallization necessary fora random texture appears at the time of hot rolling. If the randomtexture can be obtained, the plastic anisotropy is reduced and theformability is improved.

By making the Vickers hardness of the steel sheet 100 HV to 150 HV (whenC: 0.10 to 0.40% and Mn: 0.01 to 0.30%) or by making it 100 HV to 170HV, it is possible to improve the formability at the time of coldforging. If the Vickers hardness is less than 100 HV, buckling easilyoccurs during the forming at the time of cold forging and the shapedpart falls in precision, so the Vickers hardness is made 100 HV or more.Preferably, it is 110 HV or more.

If the Vickers hardness is over 170 HV, the ductility falls, buckling tooutside the plane easily occurs during thickening or other compressiondeformation, further, internal fracture easily occurs at the time ofcold forging, and the impact resistance characteristic deteriorates, sothe Vickers hardness is made 170 HV or less. To reliably secure theductility and impact resistance characteristic, the Vickers hardness ispreferably made 150 HV or less. More preferably, it is 140 HV or less.

Next, the method of production of steel sheet of the present inventionwill be explained.

The method of production of the present invention has as its basic ideato use a steel slab of the above-mentioned chemical composition andintegrally manage the hot rolling conditions and annealing conditions tocontrol the structure of the steel sheet.

First, a steel slab obtained by continuously casting molten steel of therequired chemical composition is used for hot rolling. The continuouslycast slab may be directly used for hot rolling or may be used for hotrolling after cooling once, then heating.

If cooling once, then heating the steel slab for use for hot rolling,the heating temperature is preferably 1000° C. to 1250° C. and theheating time is preferably 0.5 hour to 3 hours. If directly using thecontinuously cast steel slab for hot rolling, the temperature of thesteel slab used for the hot rolling is preferably made 1000° C. to 1250°C.

If the temperature of the steel slab or the heating temperature of thesteel slab is over 1250° C. or the heating time of the steel slab isover 3 hours, decarburization from the surface layer of the steel slabbecomes remarkable, at the time of heating before carburization andquenching, the austenite grains at the surface layer of the steel sheetabnormally grow, and the impact resistance falls. For this reason, thetemperature of the steel slab or the heating temperature of the steelslab is preferably 1250° C. or less and the heating time is preferably 3hours or less. More preferably, they are 1200° C. or less and 2.5 hoursor less.

If the temperature of the steel slab or the heating temperature of thesteel slab is less than 1000° C. or the heating time is less than 0.5hour, the microsegregation and macrosegregation occurring in castingcannot be eliminated, regions remain inside the steel slab where Si, Mn,and other alloy elements locally concentrate, and the impact resistancefalls. For this reason, the temperature of the steel slab or the heatingtemperature of the steel slab is preferably 1000° C. or more and theheating time is preferably 0.5 hour or more. More preferably they are1050° C. or more and 1 hour or more.

The finish rolling in the hot rolling is completed at 820° C. or more,preferably at 900° C. to 950° C. in temperature region. If the finishrolling temperature is less than 820° C., the steel sheet increases indeformation resistance, the rolling load remarkably rises, and, further,the amount of roll wear increases and the productivity falls. Along withthis, the recrystallization required for improving the plasticanisotropy does not sufficiently proceed, so the finish rollingtemperature is made 820° C. or more. From the viewpoint of promotingrecrystallization, it is preferably 900° C. or more.

If the finish rolling temperature is over 950° C., bulky scale formsduring passage through the run out table (ROT). Due to this scale, flawsare formed at the surface of the steel sheet. When an impact load isapplied after cold forging and carburization, quenching, and tempering,cracks easily form starting from the flaws, so the steel sheet falls inimpact resistance. For this reason, the finish rolling temperature ismade 950° C. or less. Preferably it is 920° C. or less.

When cooling the hot rolled steel sheet after finish rolling at the ROT,the cooling rate is preferably 10° C/sec to 100° C/sec. If the coolingrate is less than 10° C/sec, bulky scale is formed during the cooling.It is not possible to suppress the formation of flaws due to this andthe impact resistance falls, so the cooling rate is preferably 10° C/secor more. More preferably it is 15° C/sec or more.

If cooling from the surface layer of the steel sheet to the inside by anover 100° C/sec cooling rate, the outermost layer part is excessivelyheated and bainite, martensite, and other low temperature transformedstructures are formed. When coiling, then cooling down to 100° C. toroom temperature, then paying out the hot rolled steel sheet coil,microcracks form in the low temperature transformed structures. Themicrocracks are difficult to remove by pickling and cold rolling.

Further, if applying an impact load to the steel sheet after coldforging and carburization, quenching, and tempering, cracks advancestarting from the microcracks, so the impact resistance falls. For thisreason, to suppress the formation of bainite, martensite, and other lowtemperature transformed structures at the outermost layer part of thesteel sheet, the cooling rate is preferably 100° C/sec or less. Morepreferably it is 90° C/sec or less.

Note that, the cooling rate indicates the cooling ability received fromthe cooling facilities in a water spray section at the time when beingcooled on the ROT down to the target temperature of coiling from thetime when the hot rolled steel sheet after finish rolling is watercooled at a water spray section after passing through a non-water spraysection. It does not show the average cooling rate from the startingpoint of water spray to the temperature at which the steel sheet iscoiled up by the coiler.

The coiling temperature is made 400° C. to 550° C. This is a temperaturelower than the general coiling temperature and in particular is acondition not generally used when the content of C is high. By coilingup the hot rolled steel sheet produced under the above conditions inthis temperature range, the structure of the steel sheet can be made abainite structure comprised of fine ferrite in which carbides aredispersed.

If the coiling temperature is less than 400° C., the austenite, whichwas not transformed before coiling, transforms to hard martensite. Atthe time of paying out the hot rolled steel sheet coil, cracks form atthe surface layer of the hot rolled steel sheet and the impactresistance falls.

Furthermore, at the time of recrystallization from austenite to ferrite,since the recrystallization driving force is small, the recrystallizedferrite grains are strongly influenced in orientation by the orientationof the austenite grains and randomization of the texture becomesdifficult. For this reason, the coiling temperature is made 400° C. ormore. Preferably it is 430° C. or more.

If the coiling temperature is over 550° C., pearlite with the largelamellar spacing is formed and highly heat stable bulky needle-shapedcarbides are formed. These needle-shaped carbides remain even aftertwo-stage annealing. At the time of cold forging and otherwise formingsteel sheet, cracks are formed starting from these needle-shapedcarbides.

Further, at the time of recrystallization from austenite to ferrite,conversely the recrystallization driving force becomes too large. Inthis case as well, the result becomes recrystallized ferrite grainsheavily dependent on the orientation of the austenite grains and thetexture is not randomized. For this reason, the coiling temperature ismade 550° C. or less. Preferably it is 520° C. or less.

The hot rolled steel sheet coil is paid out and pickled, then is held intwo temperature regions for two-stage step type of annealing (two-stageannealing). By treating the hot rolled steel sheet by two-stageannealing, the stability of the carbides is controlled to promote theformation of carbides at the ferrite grain boundaries.

If cold rolling the pickled steel sheet before annealing treatment, theferrite grains are refined, so the steel sheet becomes harder to soften.For this reason, in the present invention, it is not preferable to coldroll the steel before annealing. It is preferable to perform theannealing treatment without cold rolling after the pickling.

The first stage of annealing is performed at 650 to 720° C., preferablythe A_(c1) point or less in temperature region. Due to this annealing,the carbides are coarsened and partially spheroidized and the alloyelements are made to concentrate at the carbides to thereby raise thethermal stability of the carbides.

In the first stage of annealing, the heating rate up to the annealingtemperature (below, referred to as the “first stage heating rate”) ismade 30° C/hour to 150° C/hour. If the first stage heating rate is lessthan 30° C/hour, raising the temperature takes time and the productivityfalls, so the first stage heating rate is made 3° C/hour or more.Preferably it is 10° C/hour or more.

On the other hand, if the first stage heating rate is over 150° C/hour,the temperature difference between the outer circumferential part andthe inside part of the hot rolled steel sheet coil increases, scratchesand seizing occur due to the difference in heat expansion, and reliefshapes are formed at the steel sheet surface. At the time of coldforging and other forming, cracks occur starting from the relief shapesand result in a drop in cold forgeability and a drop in formability andimpact resistance after carburizing, quenching, and tempering, so thefirst stage heating rate is made 150° C/hour or less. Preferably, it is130° C/hour or less.

The annealing temperature in the first stage of annealing (below,referred to as “the first stage annealing temperature”) is made 650° C.to 720° C. If the first stage annealing temperature is less than 650°C., the carbides become insufficient in stability and it becomesdifficult to form carbides remaining in the austenite in the secondstage of annealing. Therefore, the first stage annealing temperature ismade 650° C. or more. Preferably it is 670° C. or more.

On the other hand, if the annealing temperature exceeds 720° C., beforethe carbides rise in stability, austenite is formed and it becomesimpossible to control the above-mentioned changes in structure, so thefirst stage annealing temperature is made 720° C. or less. Preferably itis 700° C. or less.

The annealing time in the first stage of annealing (below, referred toas the “first stage annealing time”) is made 3 hours to 60 hours. If thefirst stage annealing time is less than 3 hours, the carbides becomeinsufficient in stability and it becomes difficult to form carbidesremaining in the second stage of annealing. Therefore, the first stageannealing time is made 3 hours or more. Preferably it is 5 hours ormore.

On the other hand, if the first stage annealing time exceeds 60 hours,no further stabilization of the carbides can be expected. Furthermore,the productivity drops. Therefore, the first stage annealing time ismade 60 hours or less. Preferably it is 55 hours or less.

After that, the temperature is raised to 725 to 790° C., preferably theA_(c1) point to the A₃ point in temperature region, to form austenite inthe structure. At this time, the carbides in the fine ferrite grainsdissolve in the austenite, but the carbides coarsened due to the firststage of annealing remain in the austenite.

When cooling without performing the second stage of annealing, theferrite grain size does not become larger and the ideal structure cannotbe obtained.

The heating rate up to the annealing temperature in the second stage ofannealing (below, referred to as the “second stage heating rate”) ismade 1° C/hour to 80° C/hour. At the time of the second stage ofannealing, austenite is formed and grows from the ferrite grainboundaries. At that time, by slowing the heating rate up to theannealing temperature, it becomes possible to suppress the formation ofnuclei of austenite raise the rate of coverage of the grain boundariesby carbides in the structure formed by gradual cooling after annealing.

For this reason, the second stage heating rate preferably is slow, butif less than 1° C/hour, raising the temperature takes time and theproductivity falls, so the second stage heating rate is made 1° C/houror more. Preferably it is 10° C/hour or more.

If the second stage heating rate is over 80° C/hour, the temperaturedifference between the outer circumferential part and the inside part ofthe hot rolled steel sheet coil increases, scratches and seizing occurdue to the large difference in heat expansion due to the transformation,and relief shapes are formed at the steel sheet surface. At the time ofcold forging, cracks occur starting from the relief shapes and result ina drop in cold forgeability and formability and further a drop in theimpact resistance after carburizing, quenching, and tempering as well,so the second stage heating rate is made 80° C/hour or less. Preferably,it is 70° C/hour or less.

The annealing temperature in the second stage of annealing (below,referred to as “the second stage annealing temperature”) is made 725° C.to 790° C. If the second stage annealing temperature is less than 725°C., the amount of austenite formed becomes smaller, the number ofcarbides at the ferrite grain boundaries decreases after cooling afterthe second stage of annealing, and, further, the ferrite grain sizebecomes smaller. Therefore, the second stage annealing temperature ismade 725° C. or more. Preferably it is 735° C. or more.

On the other hand, if the second stage annealing temperature exceeds790° C., it becomes difficult to make carbides remain in the austeniteand it becomes difficult to control the changes in structure, so thesecond stage annealing temperature is made 790° C. or less. Preferablyit is 770° C. or less.

The annealing time in the second stage of annealing (“second stageannealing time”) is made 3 hours to 10 hours. If the second stageannealing time is less than 3 hours, the amount of formation ofaustenite becomes small, the carbides inside the ferrite grains do notsufficiently dissolve, it becomes difficult to make the number ofcarbides at the ferrite grain boundaries increase, and, further, theferrite grain size becomes small. Therefore, the second stage annealingtime is made 3 hours or more. Preferably it is 5 hours or more.

On the other hand, if the second stage annealing time exceeds 10 hours,it becomes difficult to make carbides remain in the austenite. Further,the manufacturing costs also increase. Therefore, the second stageannealing time is made less than 10 hours. Preferably it is 8 hours orless.

After the two-stage annealing, the steel sheet is cooled by a 1° C./hour to 100° C/hour cooling rate down to 650° C.

By gradually cooling the austenite formed at the second stage ofannealing due to gradual cooling, it transforms to ferrite, carbon atomsare adsorbed at the carbides remaining in the austenite, the carbidesand austenite cover the ferrite grain boundaries, and, finally, astructure can be obtained in which a large amount of carbides arepresent at the ferrite grain boundaries. For this reason, the coolingrate is preferably slow, but if less than 1° C/hour, the time requiredfor cooling increases and the productivity falls, so the cooling rate ismade 1° C/hour or more.

Preferably it is 10° C/hour or more.

On the other hand, if the cooling rate is over 100° C/hour, theaustenite transforms to pearlite, the steel sheet increases in hardness,and the cold forgeability falls. Further, after carburization,quenching, and tempering, the impact resistance falls. Therefore, thecooling rate is made 100° C/hour or less. Preferably it is 80° C/hour orless.

Furthermore, after cooling it down to 650° C., the steel sheet is cooleddown to room temperature. The cooling rate at this time is not limited.

The atmosphere in the two-stage annealing is not particularly limited.For example, it may be any of a 95% or more nitrogen atmosphere, 95% ormore hydrogen atmosphere, or air atmosphere.

As explained above, according to the method of production of the presentinvention integrally managing the hot rolling conditions and annealingconditions and controlling the structure of the steel sheet, it ispossible to produce steel sheet excellent in formability at the time ofcold forging combining drawing and thickening and, furthermore,excellent in the hardenability required for improvement of the impactresistance after carburization, quenching, and tempering. Examples

Next, examples of the present invention will be explained, but theconditions in the examples are illustrations of conditions employed forconfirming the workability and effects of the present invention. Thepresent invention is not limited to these illustrations of conditions.The present invention can employ various conditions so long as notdeparting from the gist of the present invention and achieving theobject of the present invention.

EXAMPLE 1

A continuously cast slab (steel slab) of each of the chemicalcompositions shown in Table 1 and Table 2 (continuation of Table 1) washeated at 1240° C. for 1.8 hours, then hot rolled, cooled down to 530°C. by a 45° C/sec cooling rate on the ROT after finish hot rolling at920° C., and coiled at 520° C. to produce sheet thickness 5.2 mm hotrolled steel sheet coil.

The hot rolled steel sheet coil was paid out and pickled, then wasloaded into a box type annealing furnace. The annealing atmosphere wascontrolled to 95% hydrogen −5% nitrogen, then the coil was heated fromroom temperature to 705° C. by a 100° C/hour heating rate and was heldat 710° C. for 24 hours to obtain a uniform temperature distributioninside the hot rolled steel sheet coil.

Next, the coil was heated by a 5° C/hour heating rate to 740° C., wasfurther held at 740° C. for 5 hours, then was cooled down to 650° C. bya 10° C/hour cooling rate, then was furnace cooled down to roomtemperature to prepare a sample for evaluation of performance. Thestructure of the sample was observed by the method explained above andthe ferrite grain size and number of carbides were measured.

TABLE 1 No C Si Mn P S Al N O Ti Cr Mo B Nb V Cu Remarks 1 0.35 0.210.53 0.0105 0.0027 0.065 0.0067 0.01 0.041 0.0300 0.1760 0.0090 Inv.steel 2 0.24 0.22 0.59 0.0075 0.0040 0.066 0.0021 0.01 0.035 0.45000.1330 0.0055 Inv. steel 3 0.18 0.03 0.41 0.0165 0.0026 0.029 0.00500.01 0.094 0.1500 0.4390 0.0022 Inv. steel 4 0.31 0.13 0.31 0.00030.0053 0.072 0.0048 0.02 0.050 0.1300 0.2510 0.0028 Inv. steel 5 0.250.09 0.39 0.0021 0.0088 0.055 0.0079 0.01 0.085 0.0500 0.3580 0.0098Inv. steel 6 0.13 0.04 0.7 0.0088 0.0066 0.072 0.0031 0.02 0.058 0.49000.0170 0.0018 Inv. steel 7 0.14 0.04 0.32 0.0133 0.0081 0.052 0.00170.02 0.094 0.3700 0.3190 0.0093 Inv. steel 8 0.21 0.03 0.93 0.01470.0077 0.074 0.0021 0.01 0.085 0.3300 0.1440 0.0042 Inv. steel 9 0.290.22 0.4 0.0118 0.0082 0.036 0.0071 0.01 0.040 0.2700 0.0170 0.0082 Inv.steel 10 0.31 0.23 0.6 0.0040 0.0068 0.037 0.0010 0.01 0.039 0.24000.4480 0.0020 Inv. steel 11 1.20 0.02 0.78 0.0176 0.0049 0.073 0.00920.00 0.020 0.1600 0.0070 0.0024 Comp. steel 12 0.27 1.50 1.0 0.01700.0029 0.013 0.0004 0.00 0.037 0.2900 0.2990 0.0096 Comp. steel 13 0.370.21 3.3 0.0189 0.0052 0.046 0.0099 0.01 0.027 0.1400 0.2770 0.0077Comp. steel 14 0.39 0.27 0.83 0.0091 0.0059 0.018 0.0023 0.01 0.0381.1000 0.4070 0.0099 Comp. steel 15 0.35 0.26 0.37 0.0143 0.0031 0.0180.0063 0.00 0.028 0.0200 0.1340 0.0000 Comp. steel 16 0.22 0.05 0.360.0109 0.0032 0.067 0.0008 0.0062 0.001 0.18 0.081 0.0014 0.028 0.083Inv. steel 17 0.35 0.2 0.47 0.0137 0.0004 0.052 0.0013 0.0127 0.054 0.450.241 0.0057 Inv. steel 18 0.37 0.07 0.95 0.0076 0.0008 0.057 0.00980.0149 0.097 0.23 0.074 0.0019 Inv. steel 19 0.37 0.23 0.76 0.00860.0058 0.073 0.0057 0.0076 0.076 0.18 0.149 0.0007 Inv. steel 20 0.320.17 0.61 0.0197 0.0024 0.011 0.0009 0.0153 0.083 0.27 0.295 0.0019 Inv.steel 21 0.22 0.27 0.57 0.0059 0.0013 0.048 0.0041 0.0064 0.005 0.330.438 0.0085 Inv. steel 22 0.12 0.19 0.32 0.0135 0.0032 0.078 0.00960.0107 0.083 0.06 0.167 0.0051 0.024 Inv. steel 23 0.25 0.13 0.8 0.00560.0046 0.087 0.0064 0.0001 0.04 0.32 0.33 0.0079 Inv. steel 24 0.31 0.160.41 0.0142 0.0097 0.026 0.0054 0.0081 0.003 0.14 0.473 0.0049 Inv.steel 25 0.22 0.23 0.42 0.0052 0.003 0.093 0.0055 0.0197 0.05 0.21 0.3490.0081 Inv. steel 26 0.34 0.24 0.68 0.01 0.0081 0.024 0.0077 0.00920.085 0.05 0.212 0.0024 0.044 0.09  Comp. steel 27 0.11 0.09 0.45 0.01240.0068 0.008 0.0059 0.0151 0.044 0.41 0.109 0.0074 0.21  Comp. steel 280.2 0.28 0.5 0.0196 0.0007 0.003 0.0060 0.0137 0.012 0.3 0.406 0.0013Comp. steel 29 0.34 0.1 0.71 0.0004 0.003 0.014 0.0059 0.0093 0.092 0.020.189 0.0082 0.22  0.004 Comp. steel 30 0.37 0.09 0.8 0.007 0.0071 0.0430.0032 0.0023 0.016 0.33 0.253 0.0044 0.36  Comp. steel 31 0.33 0.210.53 0.0104 0.0028 0.066 0.0069 0.01 0.039 0.0320 0.1730 0.0093 Inv.steel 32 0.36 0.04 0.32 0.011 0.0032 0.067 0.0008 0.0062 0.001 0.180.081 0.0014 0.054 0.087 Inv. steel 33 0.34 0.20 1.21 0.0105 0.00270.065 0.0075 0.01 0.041 0.0320 0.1720 0.0095 Inv. steel 34 0.22 0.051.83 0.01 0.0032 0.066 0.0007 0.0062 0.001 0.18 0.081 0.0014 0.032 0.096Inv. steel 35 0.37 0.21 1.3 0.0104 0.0029 0.065 0.0063 0.01 0.038 0.02800.1690 0.0096 Inv. steel 36 0.22 0.06 2.19 0.0109 0.0035 0.067 0.00080.0062 0.001 0.18 0.081 0.0014 0.036 0.065 Inv. steel

TABLE 2 (Continuation of Table 1) (mass %) No W Ta Ni Sn Sb As Mg Ca YZr La Ce Remarks 1 Inv. steel 2 Inv. steel 3 Inv. steel 4 Inv. steel 5Inv. steel 6 Inv. steel 7 Inv. steel 8 Inv. steel 9 Inv. steel 10 Inv.steel 11 Comp. steel 12 Comp. steel 13 Comp. steel 14 Comp. steel 15Comp. steel 16 Inv. steel 17 0.033 0.043 Inv. steel 18 0.091 Inv. steel19 0.014 Inv. steel 20 0.032 0.022 Inv. steel 21 0.0424 Inv. steel 220.031 0.046 Inv. steel 23 0.046 Inv. steel 24 0.017 0.024 Inv. steel 250.027 Inv. steel 26 0.15 Comp. steel 27 0.049 0.05 Comp. steel 28 0.330.012 0.041 Comp. steel 29 0.024 Comp. steel 30 0.002 Comp. steel 31Inv. steel 32 Inv. steel 33 Inv. steel 34 Inv. steel 35 Inv. steel 36Inv. steel

Table 3 shows the ferrite grain size (μm), average carbide grain size(μm), pearlite area ratio (%), Vickers hardness (HV), number of grainboundary carbides/number of grain carbides, X-ray intensity ratio I1/I0,“r” value anisotropy index |Δr|, and critical cooling rate (° C/sec)shown in Table 1 and Table 2. If I1/I0 is 1 or more, therecrystallization in hot rolling does not sufficiently proceed and thesteel sheet becomes larger in plastic anisotropy. Note that, the “r”value anisotropy index |Δr| was found by a tensile test

TABLE 3 Average No. of grain Ferrite carbide Pearlite boundary Criticalgrain area grain Vickers carbides/No. cooling size size rate hardness ofgrain rate No (μm) (μm) (%) (HV) carbides I1/I0 |Δr| (° C./sec) Remarks1 17 0.7 1.7 120 5.66 0.75 0.17 29.8 Inv. ex. 2 23 1.1 1.2 118 5.85 0.810.19 29.8 Inv. ex. 3 16 0.8 0.8 105 3.88 0.80 0.18 29.9 Inv. ex. 4 131.1 1.3 109 5.54 0.76 0.17 30.0 Inv. ex. 5 20 0.9 1.0 108 3.48 0.69 0.1529.9 Inv. ex. 6 12 1.1 0.0 115 7.11 0.67 0.14 29.9 Inv. ex. 7 13 1.3 1.0100 6.11 0.64 0.13 30.1 Inv. ex. 8 11 1.2 1.3 124 4.23 0.64 0.13 29.7Inv. ex. 9 20 1.1 1.2 113 6.07 0.79 0.18 29.9 Inv. ex. 10 17 1.4 1.9 1226.59 0.65 0.13 29.8 Inv. ex. 11 18 1.1 9.1 167 5.36 0.69 0.15 30.0 Comp.ex. 12 10 1.0 1.5 154 6.86 0.69 0.15 29.4 Comp. ex. 13 14 1.3 12.3 1785.69 0.72 0.16 13.2 Comp. ex. 14 20 1.0 1.2 133 0.91 0.78 0.18 29.6Comp. ex. 15 12 1.1 1.9 114 6.04 0.79 0.18 311.0 Comp. ex. 16 11 0.8 1.8105 5.02 0.63 0.13 30.1 Inv. ex. 17 26 1.0 1.6 119 4.19 0.66 0.14 29.8Inv. ex. 18 13 1.2 1.0 134 5.05 0.78 0.18 29.6 Inv. ex. 19 11 1.2 1.8130 5.36 0.75 0.17 29.6 Inv. ex. 20 20 1.1 1.0 122 4.17 0.65 0.13 29.7Inv. ex. 21 11 1.1 1.0 116 5.57 0.74 0.16 29.9 Inv. ex. 22 17 1.1 0.6100 6.17 0.82 0.19 30.1 Inv. ex. 23 14 1.2 1.0 122 6.19 0.77 0.17 29.7Inv. ex. 24 12 1.1 1.8 113 7.12 0.81 0.19 29.9 Inv. ex. 25 17 0.8 1.4111 5.38 0.74 0.16 29.9 Inv. ex. 26 13 1.2 1.4 126 6.46 0.80 0.18 29.6Comp. ex. 27 19 1.1 1.2 105 6.93 0.74 0.16 30.0 Comp. ex. 28 23 0.9 1.8113 6.58 0.70 0.15 29.9 Comp. ex. 29 18 1.2 0.5 125 4.62 0.71 0.15 29.7Comp. ex. 30 14 1.1 0.7 129 4.93 0.68 0.14 29.7 Comp. ex. 31 18 1.1 1.5125 5.72 0.75 0.17 29.3 Inv. ex. 32 13 0.8 0.2 109 5.17 0.63 0.13 29.7Inv. ex. 33 18 1.3 1.2 125 5.60 0.75 0.16 31.2 Inv. ex. 34 13 1.4 1.9109 4.97 0.63 0.13 30.9 Inv. ex. 35 18 1.3 1.4 128 5.75 0.75 0.17 29.5Inv. ex. 36 13 1.4 1.0 110 5.23 0.63 0.11 29.8 Inv. ex.

In general, if the anisotropy index |Δr| obtained from the “r” values inparallel to the sheet surface and in three directions is over 0.2, thedrawability falls. Therefore, to secure excellent formability, a |Δr|not over 2 is demanded.

The critical cooling rate was found by preparing a CCT graph. If coolinghot rolled steel sheet by a cooling rate slower than the found criticalcooling rate, the hardenability at the time of hardening after forming apart becomes poorer and pearlite structures are formed so sufficientstrength cannot be obtained. For this reason, the critical cooling ratemust be small in order to obtain a high hardening strength. If thecritical cooling rate is 280° C/sec, it can be judged that thehardenability is improved.

In the invention examples shown in Table 3, the average carbide grainsize is 0.4 to 2.0 μm, the pearlite area ratio is 6% or less, the numberof grain boundary carbides/number of grain carbides is over 1, and theI1/I0 is less than 1, so the Vickers hardness is 100 HV to 170 HV inrange and |Δr| is less than 0.2. In the comparative examples using thecomparative steel sheets, the Vickers hardness is over 150, while thenumber of grain boundary carbides/number of grain carbides becomes lessthan 1. In the comparative steel sheet in which B is not added (inTables 1 and 2, No. 15), the critical cooling rate is over 280° C/secand the hardenability falls.

EXAMPLE 2

A method of production of conditions outside the scope of conditionsprescribed in the present invention was applied to the 12 types of steelof the Invention Steel Nos. 1 to 5, Nos. 16 to 19, Nos. 31, No. 33, andNo. 35. Table 4 shows the manufacturing conditions, while Table 5 showsthe ferrite grain size (μm), Vickers hardness (HV), number of grainboundary carbides/number of grain carbides, X-ray intensity ratio I1/I0,“r” value anisotropy index |Δr|, and critical cooling rate (° C/sec) ofsteel sheets produced under the manufacturing conditions shown in Table4.

TABLE 4 Hot rolling conditions Annealing conditions Finish 1st stage 2ndstage Cooling rolling Coiling Heating Holding Holding Heating HoldingHolding rate down temp. temp. rate temp. time rate temp. time to 650° C.No (° C.) (° C.) (° C./hour) (° C.) (hours) (° C./hour) (° C.) (hours)(° C./hour) 1 720 520 60 710 24 60 740 5 80 2 970 520 60 710 24 60 740 580 3 920 350 60 710 24 60 740 5 80 4 920 570 60 710 24 60 740 5 80 5 920520 60 600 24 60 740 5 80 16 920 520 60 710 2 60 740 5 80 17 920 520 60710 24 60 720 5 80 18 920 520 60 710 24 60 820 5 80 19 920 520 60 710 2460 740 18 80 31 720 520 60 710 24 60 740 5 80 33 920 520 60 600 24 60740 5 80 35 920 520 60 710 24 60 820 5 80

TABLE 5 Average No. of grain Ferrite carbide Pearlite boundary Criticalgrain grain area Vickers carbides/No. cooling size size rate hardness ofgrain rate No (μm) (μm) (%) (HV) carbides I1/I0 |Δr| (° C./sec) Remarks1 17 0.9 1.3 120 5.66 1.2 0.33 29.8 Comp. ex. 2 23 1.1 2.1 118 5.85 1.60.47 29.8 Comp. ex. 3 16 0.8 0.8 105 3.88 1.1 0.29 29.9 Comp. ex. 4 131.1 3.1 109 5.54 1.3 0.36 30.0 Comp. ex. 5 20 0.9 4.5 108 0.81 0.82 0.1929.9 Comp. ex. 16 11 1.1 3.8 105 0.53 0.79 0.18 30.1 Comp. ex. 17 26 1.01.2 119 0.92 0.73 0.16 29.8 Comp. ex. 18 13 2.4 9.4 161 0.67 0.68 0.1429.6 Comp. ex. 19 11 1.9 10.2 155 0.83 0.62 0.12 29.6 Comp. ex. 31 170.9 0.4 125 5.72 1.2 0.33 29.8 Comp. ex. 33 18 1.3 4.8 130 0.92 0.750.17 31.2 Comp. ex. 35 18 1.7 8.2 163 0.83 0.75 0.17 29.5 Comp. ex.

It will be understood that making the finish rolling temperature in hotrolling or the coiling temperature a temperature outside of the scope ofconditions prescribed in the present invention invites a drop in therecrystallization and has a large effect on the randomization of thetexture and as a result causes the value of |Δr| to rise. Further, itwill be understood that if making the annealing conditions outside thescope of conditions prescribed in the present invention, the number ofgrain boundary carbides/number of grain boundary carbides becomes 1 orless and the state of distribution of carbides greatly changes.

INDUSTRIAL APPLICABILITY

As explained above, according to the present invention, it is possibleto provide steel sheet excellent in hardenability and formability as amaterial and a method of production of the same. The steel sheet of thepresent invention is suitable for forming a part by cold forging such asthickening to obtain a gear or other part. Accordingly, the presentinvention has high applicability in the manufacture of steel sheet andindustries utilizing it.

1. A steel sheet consisting of, by mass %, C: 0.10 to 0.70%, Si: 0.01 to0.30%, Mn: 0.30 to 3.00%, Al: 0.001 to 0.10%, Cr: 0.010 to 0.50%, Mo:0.0010 to 0.50%, B: 0.0004 to 0.01%, Ti: 0.001 to 0.10%, P: 0.02% orless, S: 0.01% or less, N: 0.0200% or less, O: 0.0200% or less, Sn:0.05% or less, Sb: 0.05% or less, As: 0.05% or less, Nb: 0.10% or less,V: 0.10% or less, Cu: 0.10% or less, W: 0.10% or less, Ta: 0.10% orless, Ni: 0.10% or less, Mg: 0.05% or less, Ca: 0.05% or less, Y: 0.05%or less, Zr: 0.05% or less, La: 0.05% or less, and Ce: 0.05% or less anda balance of Fe and unavoidable impurities, wherein a metal structure ofthe steel sheet includes carbide having an average grain size of 0.4 μmto 2.0 μm, perlite having an area ratio of 6% or less and ferritewherein a ratio of a number of the carbides at ferrite grain boundariesto a number of the carbides in ferrite grains of over 1; and I1/I0<1 issatisfied when an X-ray diffraction intensity at {211}<011>at a plane ofa part of ½ sheet thickness of the steel sheet is denoted by “I1” and anX-ray diffraction intensity at {100}<011>is denoted by “I0”, the steelsheet having a Vickers hardness of 100 HV to 150 HV.
 2. A method ofproduction for producing steel sheet according to claim 1 comprising hotrolling a steel slab of a chemical composition according to claim 1 withfinish rolling temperature between 820° C. and 950° C., to obtain hotrolled steel sheet; coiling the hot rolled steel sheet at 400° C. to550° C.; pickling the coiled hot rolled steel sheet; heating the pickledhot rolled steel sheet to an annealing temperature of 650° C. to 720° C.by a heating rate of 30° C/hour to 150° C/hour and holding the steelsheet for 3 hours to 60 hours as a first stage of annealing; next,heating the hot rolled steel sheet to an annealing temperature of 725°C. to 790° C. by a heating rate of 1° C/hour to 80° C/hour and holdingthe steel sheet for 3 hours to less than 10 hours as a second stage ofannealing; and, next, cooling the annealed hot rolled steel sheet to650° C. by a cooling rate of 1° C/hour to 100° C/hour.